Effects of Mn on Isothermal Transformation Microstructure and Tensile Properties in Medium- and High-carbon Steels

Abstract

The effects of Mn addition on the microstructure formed through an isothermal transformation at 873 K and its tensile properties were investigated over a wide range of concentrations in medium- and high-carbon steels with 0.4–1.0 mass% C. The Mn addition changed the pearlite transformation mode from relatively slow non-partitioning pearlite with a small amount of proeutectoid ferrite to an extremely slow partitioning pearlite without any proeutectoid ferrite in the hypoeutectoid steel. The transformation rate of the partitioning pearlite associated with the Mn partitioning between ferrite and cementite in Fe–C–M alloy was more than three orders of magnitude lower than the pearlite transformation in Fe–C binary alloy at a given interlamellar spacing. The decrease in the transformation rate in the non-partitioning and partitioning pearlite transformation was caused by the decrease in the carbon flux controlling the pearlite transformation, which can be explained by the theory of local equilibrium at the austenite/ferrite and austenite/cementite interphases. The Mn addition increased the thermal stability of the lamellar cementite. Corresponding to the change in the transformation microstructure, the Mn addition improved the tensile properties in the pearlite steel, particularly the strength and local deformability balance, regardless of the difference in the transformation mode between the non-partitioning and the partitioning transformation, unless the proeutectoid cementite precipitated at prior austenite grain boundaries. The strength increase of the pearlite after the Mn addition was caused by the refinement of the interlamellar spacing and/or the increase of the lattice strain in the pearlitic ferrite.

1. Introduction

Medium- (0.3–0.8%C) and high-carbon (0.8–2.0%C) steels, which are widely used in many applications as tool steel, have a ferrite-pearlite duplex or fully pearlite microstructure at the non-heat-treated state (% = mass%). The mechanical properties of this type of steel strongly depend on the microstructural characteristics of pearlite, e.g., the fraction and interlamellar spacing of the pearlite. Therefore, they are optimized precisely by controlling C content and the transformation temperature. Additionally, the alloy elements are important for the microstructural control of the pearlite in medium- and high-carbon steels. From the viewpoint of thermodynamics, Mn addition moves a eutectoid point with a lower C region and a lower temperature, which leads to the reduction of the proeutectoid ferrite fraction.¹⁾ Additionally, Cr addition increases the eutectoid temperature, which contributes to the high strength and strain hardenability of the pearlite, owing to the refinement of the interlamellar spacing.²⁾ Moreover, it has been reported that Si addition suppresses the spheroidization of the cementite lamellae.^3,4) Their alloy effects enable the industrial production of an ultra-high strength steel wire, i.e., the steel code, at low cost.^{2,3,4,5,6,7,8)} However, in conventional steels, the concentration of the alloy elements is generally limited to less than approximately 1.0%. Thus, the effect of the alloy addition on the microstructure and the mechanical properties of the medium- and high-carbon steels in the higher concentration region are unclear. In particular, we should pay attention to the fact that the behavior of the pearlite during the transformation changes drastically with a high concentration of alloy addition. Ridley et al. have used a eutectoid steel containing approximately 1.0%Mn and reported that the Mn is partitioned between the ferrite and the cementite in the pearlite transformation, when the transformation temperature is relatively high, which leads to a significant retardation of the pearlite transformation rate.^9,10) The slow pearlite transformation accompanied with alloy partitioning between ferrite and cementite has also been reported in Cr added to eutectoid steel by Chance and Ridley.¹¹⁾ Hillert et al. proposed that the local equilibrium (LE) theory can explain the alloy partitioning behavior in the pearlite transformation.¹²⁾ Indeed, Hutchinson et al. conducted experiments and reported that the pearlite transformation can be classified into two types: one type is controlled by fast C diffusion (non-partitioning pearlite), while the other is controlled by a very slow Mn diffusion (partitioning pearlite), and that the classification between the partitioning and non-partitioning pearlite is based on the LE theory with regard to the Fe–C–Mn alloy.¹³⁾ Therefore, to fully understand the effect of the alloy addition on the microstructure and mechanical properties of the pearlite, the microstructural characteristics of the pearlite should be investigated by distinguishing between the non-partitioning and partitioning pearlite. In this study, medium- and high-carbon steels with a high concentration of Mn up to 3% were prepared and the effect of the Mn addition on the isothermal transformation microstructure, and the tensile properties of the Fe–C–Mn alloys were investigated. The Mn partitioning among the austenite, ferrite, and cementite phases in the pearlite transformation are discussed on the basis of LE theory, and the structure-property relationship is clarified to understand the mechanism of pearlite strengthening resulting from the Mn addition.

2. Experimental Procedures

Various medium- and high-carbon steels with different C and Mn balance were prepared with the chemical compositions listed in Table 1. In this study, each steel was named according to the Mn–C balance; for example, 40C3Mn (0.40%C-3.0%Mn steel). The ingots fabricated by a vacuum melting furnace were sufficiently homogenized at 1523 K for 36 ks, and then hot forged to a bar shape with a diameter of 8.0 mm. Each specimen cut from the bar samples was austenitized at 1293 K for 1.8 ks and directly subjected to isothermal holding at 873 K for various periods to decompose the austenite, followed by water cooling. Figure 1 shows the isothermal phase diagram of Fe–C–Mn system at 873 K, which was calculated using Thermo-Calc software with TCFE9 database. Here, u-fraction is defined as the site fraction of an element with reference to the substitutional sublattices only. Therefore, u-fraction of Mn means atomic fraction of Mn against the sum of Fe and Mn. In addition, the extended A_e3 and A_cm lines are represented by the dashed lines, while the boundaries between the partitioning and non-partitioning LE (PLE/NPLE) for austenite to ferrite (γ-α) and austenite to cementite (γ-θ) formations are indicated by the dashed-dotted lines. This diagram shows that all steels were in the ferrite and cementite two-phase region as an equilibrium state at 873 K. However, the 0.4%C series, 40C0Mn, 40C2Mn, and 40C3Mn, were not in the coupled zone, where austenite was supersaturated with respect to both ferrite and cementite. This suggests the preferential precipitation of the proeutectoid ferrite before the eutectoid transformation. The microstructure was characterized by means of optical microscopy (OM) and scanning electron microscopy (SEM). The cross-sectional surface at the central part of the heat-treated sample was mechanically gridded with emery papers (#220–#4000) and thereafter polished with colloidal silica. Finally, the surface was etched with 5% nital solution for microstructural observation. The interlamellar spacing of the pearlite was evaluated in various pearlite colonies, where the cementite lamellae lied perpendicular to the observation surface. Moreover, to evaluate the growth rate of the pearlite, the diameter of the largest pearlite nodule was measured using SEM in the isothermally held specimens. It is well known that pearlite growth occurs under steady state. Therefore, assuming that the largest pearlite nodule is the nodule that forms at the very beginning of the pearlite transformation, the change ratio of the nodule diameter against holding time corresponds to the pearlite growth rate.^9,13,14) The hardness was evaluated using a Vickers hardness testing machine (MMT-X, developed by Matsuzawa Co., Ltd.) with a load of 9.8 N. The measurements were performed for 10 times and the mean value was calculated. The Mn partitioning amongst ferrite, cementite, and austenite phases was accurately measured using a spherical aberration corrected scanning transmission electron microscope with energy-dispersive-spectroscopy (Cs-STEM/EDX, JEM-ARM200F developed by JEOL). The tensile testing was carried out using an Instron type testing machine (AG-100kNG·RX, developed by Shimazu Corp.) at an initial strain rate of 1.7×10⁻³ s⁻¹ for round-bar test pieces with a gauge diameter of φ3.0 mm × 10 mm. To quantify the lattice strain (micro-strain) in the ferrite phase ε_L, the full width at a half maximum β was measured by means of X-ray diffraction analysis (Smartlab-BBex, developed by Rigaku Corp.) at each diffraction angle 2θ after the separation of Kα1 and Kα2, and was plotted according to the following Williamson-Hall equation.¹⁵⁾   

$$\beta \frac{cos\theta }{{\lambda }_W}=\frac{0.9}{D}+2{\epsilon }_L\frac{sin\theta }{{\lambda }_W}$$(1)

where D and λ_W denote the crystallite size and X-ray wavelength (Cu-Kα, 0.15418 nm), respectively. In this analysis, the (200) diffraction was deselected owing to the high anisotropy, while the artifact introduced by the equipment was removed using β for a fully annealed interstitial atom free (IF) steel (Fe-0.0004%C-0.024%Ti alloy). Recently, the Williamson-Hall equation was developed to characterize the dislocations, i.e., the modified Williamson-Hall/modified Warren-Averbach methods. However, this study simply adopted the normal method expressed by Eq. (1), because, in the modified method, the lattice strain is converted into the dislocation density.

Table 1. Chemical composition of medium- and high-carbon steels used in this study.

	C	Si	Mn	P	S	Al	N	Fe
40C0Mn	0.41	<0.02	<0.02	<0.005	<0.0005	<0.005	<0.0005	bal.
40C2Mn	0.41	<0.02	1.97	<0.005	<0.0005	<0.005	0.0008	bal.
40C3Mn	0.39	<0.02	3.01	<0.005	<0.0005	<0.005	0.0015	bal.
50C2Mn	0.50	<0.02	2.00	<0.005	<0.0005	<0.005	0.0012	bal.
60C2Mn	0.60	<0.02	2.01	<0.005	<0.0005	<0.005	0.0018	bal.
100C0Mn	1.03	<0.02	<0.02	<0.005	<0.0005	<0.005	0.0014	bal.
100C2Mn	1.02	<0.02	2.02	<0.005	<0.0005	<0.005	0.0012	bal.

Fig. 1.

Isothermal phase diagram of Fe–C–Mn system at 873 K with bulk compositions of used steels and PLE/NPLE boundaries.

3. Results and Discussion

3.1. Effect of Mn Addition on Microstructure Formed via Pearlite Transformation

3.1.1. Effect of Mn Addition on Pearlite Transformation Behavior

The isothermally transformed microstructure mainly depends on C content. As may be expected from the isothermal phase diagram shown in Fig. 1, steels with C content higher than 0.5% did not exhibit proeutectoid ferrite formation. However, the precipitation of the proeutectoid cementite could not be suppressed upon cooling from the austenitization temperature to the isothermal holding temperature as C content increased. Consequently, the film shaped cementite decorated the prior austenite grain boundaries in the 1.0%C steels, 100C0Mn, and 100C2Mn, as shown in Fig. 2, for example. The film cementite became wider as Mn content increased, because the Mn addition decreased the eutectoid C content.¹⁾ Moreover, the effect of the Mn addition on the transformed microstructure was more obvious in the hypoeutectoid 0.4%C steels.

Fig. 2.

SEM microstructure images of (a) 50C2Mn and (b) 100C2Mn isothermally held at 873 K for 600 s after austenitization.

Figure 3 shows the OM and SEM images of (a, b) 40C0Mn, (c, d) 40C2Mn, and (e, f) 40C3Mn, which were isothermally held at 873 K for various times. Austenite completely decomposed to duplex microstructure consisting of the proeutectoid ferrite indicated by the white arrows, and the pearlite in (a) 40C0Mn and (c) 40C2Mn. However, there existed a difference in the morphology of the proeutectoid ferrite and the pearlite fraction; namely, F_P between them. It is particularly interesting that F_P markedly increased by the 2%Mn addition. Consequently, 40C2Mn had almost fully pearlite, although it is a hypoeutectoid steel. Moreover, (e) 40C3Mn exhibited a clearly different transformation behavior in comparison with these two steels. The pearlite (P) formed directly at the prior austenite grain boundaries without any proeutectoid ferrite, and more than half of austenite, which transformed to martensite (M) after quenching, remained even after holding for 3.6 ks. This clearly demonstrated that the pearlite transformation was remarkably retarded by Mn addition, while the transformation kinetics discontinuously decreased when the Mn content increased from 2% to 3%. Actually, the pearlite growth rate (v_P) was measured at 21.8, 1.08, and 0.03 μm/s in 40C0Mn, 40C2Mn, and 40C3Mn, respectively. By examining the eutectoid transformation product in each alloy, it was observed that the lamellar structure with approximately 150 nm in interlamellar spacing (λ) was well developed in (d) 40C2Mn and (f) 40C3Mn. However, the lamellar cementite was broken out and spheroidized in (b) 40C0Mn. It is known that cementite with a spherical shape is densely distributed in pearlite transformed at a relatively lower temperature, owing to the non-cooperative growth of the two constitutive phases. This type of eutectoid transformation product is typically considered as degenerated pearlite.¹⁶⁾ However, it was confirmed that cementite had a lamellar structure immediately after the pearlite transformation (Fig. 4). This suggests that the lamellar pearlite developed during the eutectoid transformation, but became spheroidized for a short period during the subsequent holding in 40C0Mn. From these results, it can be concluded that Mn addition in medium-C steel exerts the following main effects on the transformation microstructure: i) inhibition of proeutectoid ferrite formation; ii) reduction of pearlite growth rate; iii) thermal stabilization of lamellar cementite. Next, we will discuss these effects in the order stated.

Fig. 3.

OM and SEM microstructure images of (a, b) 40C0Mn (600 s), (c, d) 40C2Mn (600 s), and (e, f) 40C3Mn (3.6 ks) isothermally held at 873 K.

Fig. 4.

SEM image showing cooperative growth of ferrite and cementite on eutectoid transformation in 40C0Mn steel isothermally transformed at 873 K for 10 s.

3.1.2. Inhibition of Proeutectoid Ferrite Formation

Figure 5 is a schematic illustration showing (a) the change in the average C concentration in the untransformed austenite by ferrite formation in Fe–C binary phase diagram, and (b) the corresponding one-dimensional C concentration profile across the growing interface between the proeutectoid ferrite and the untransformed austenite, when a hypoeutectoid steel is held below the eutectoid temperature. Assuming that LE is satisfied at the ferrite/austenite growing interface, C concentration at the interface in the austenite region corresponds to C^γ/α and is locally higher than C^γ/θ, as indicated by the gray color. This means that if cementite precipitates at the interface, pearlite transformation can continue even when the average C concentration in the untransformed austenite (C₀) is lower than C^γ/θ.¹²⁾ In fact, Liu et al. observed the formation of ferrite-pearlite with F_P larger than the value expected from C^γ/θ and C₀ in 0.3%C-2.0%Mn steel, and reported that F_P becomes larger with the coarsening of prior austenite grain size.¹⁷⁾ Additionally, the increase in F_P related with prior austenite coarsening was confirmed in 0.4%C-1.2%Mn steel.¹⁸⁾ From these experimental results, it is thought that a transition from ferrite to pearlite transformation occurs when the velocity of the ferrite/austenite growing interface becomes sufficiently slow for cementite to precipitate at the interface. According to this theory, the increase in F_P induced by the Mn addition can be attributed to the decrease in the velocity at the ferrite/austenite interface, i.e., solute-drag effect.^19,20) In particular, when the proeutectoid ferrite grows under PLE mode, the transition from ferrite to pearlite transformation occurs before the proeutectoid ferrite grows sufficiently, and does not result in an obvious proeutectoid ferrite formation.

Fig. 5.

Schematic illustration of (a) change in average carbon content in untransformed austenite by formation of proeutectoid ferrite, and (b) the corresponding carbon profile at the moving interface between proeutectoid ferrite and untransformed austenite during ferrite-pearlite transformation in hypoeutectoid steel.

3.1.3. Understanding of Pearlite Transformation Rate Using Local Equilibrium Theory

The pearlite transformation occurs in a steady-state condition by the cooperative growth of ferrite and cementite at the growth interface, which is caused by C diffusion in the vicinity of the pearlite/austenite interface, as shown in Fig. 6. Therefore, the pearlite growth rate (v_P) discussed on the basis of C diffusion-controlled process, which is called Zener-Hillert model.^21,22) According to their theory, the relationship between v_P and λ is expressed by the following equation.   

$$v_{\text{P}}{\lambda }^{\text{n}}=\text{constant}$$(2)

Here, n is a positive number between 2 and 3. When the pearlite transformation is solely controlled by volume diffusion in austenite or pearlite/austenite interfacial diffusion of C, n is equal to 2 or 3, respectively. Figure 7 shows the change in v_P as a function of λ in 0.80%C,¹⁴⁾ 0.8%C-1.0%Mn,⁹⁾ and 0.54%C-3.51%Mn¹³⁾ steels. Moreover, 0.80%C and 0.8%C-1.0%Mn steels exhibit a good linear relationship between v_P and λ in a double-logarithmic graph, which reveals that Eq. (2) is established in both materials. Each n value was evaluated as 2.38 in the 0.80%C steel, and as 2.42 in the 0.8%C-1.0%Mn steel, which suggests that the pearlite transformation is controlled by both the volume diffusion in austenite and the interface diffusion of C. Indeed, a recent study using phase field simulation suggested the multi-pass diffusion of C in pearlite transformation.²³⁾ Assuming that 0.54%C-3.51%Mn steel has the same n value as other steels, as indicated by the dashed line, it was found that the data obtained from 40C0Mn, 40C2Mn, and 40C3Mn (●) in this study are in agreement with the data obtained by previous studies. This figure reveals that the pearlite growth rate can be classified into three levels, which are at least dependent on the Mn content at a given λ: approximately 40 μm/s (without Mn, 40C0Mn), 4.0 μm/s (low Mn, 40C2Mn), and 0.1 μm/s (high Mn, 40C3Mn) at 100 nm in λ. As shown in Fig. 6, the pearlite transformation was controlled by the flux of the C diffusion (J_C) and is expressed by the following equation with C concentration (C_C) and mobility of the C diffusion including the volume and interfacial diffusions (M) in terms of the chemical potential gradient of C (μ_C).   

$$J_C=-M{C}_C\frac{\partial {\mu }_C}{\partial x}\approx -M{C}_C\frac{RTln\left(\raisebox{1ex}{$a_C^{\gamma /\alpha }$}\!\left/ \!\raisebox{-1ex}{$a_C^{\gamma /\theta }$}\right.\right)}{\raisebox{1ex}{$\lambda $}\!\left/ \!\raisebox{-1ex}{$2$}\right.}$$(3)

Here, x is the distance along the axis parallel to the pearlite/austenite interface. Additionally, a_C^i/ii denotes the C activity in the i phase at the interface between the i and ii phases. The most important consideration in Eq. (3) is the difference in the C activity between a_C^γ/α and a_C^γ/θ controls v_P. If a_C^γ/θ is higher than a_C^γ/α, the pearlite transformation does not occur. Therefore, to understand the effect of the Mn addition on the transformation kinetics of pearlite, we will discuss the C activity at the hypoeutectoid steel interfaces with three different Mn concentrations on the basis of LE theory. Figure 8 shows the isothermal phase diagram of Fe–C–Mn system below the eutectoid temperature. The Mn content of the bulk composition (○) was different amongst (a) no-Mn, (b) low-Mn, and (c) high-Mn. In the case of binary Fe–C alloy (a), ferrite preferred to form in austenite matrix while the ferrite/austenite interface maintained the unique tie-line A₁B₁. Subsequently, when cementite precipitated in the C-supersaturated austenite, the tie-line of the cementite/austenite interface was C₁D₁. The iso-C activity line passing through B₁ and C₁ is indicated by a_C,1^γ/α and a_C,1^γ/θ with dashed lines, respectively. Thus, the C flux controlled the pearlite transformation in binary Fe–C alloy, J_C,₁ is uniquely given by the difference between a_C,1^γ/α and a_C,1^γ/θ. In the case of the low-Mn steel, whose Mn concentration was lower than the PLE/NPLE boundary for γ-α (b), the ferrite formation occurred under NPLE mode. In this case, because the formed ferrite had the same Mn concentration as the bulk composition, the operating tie-line at the ferrite/austenite interface was drawn by A₂B₂. Given that the Mn-spike at the growing NPLE ferrite/austenite interface exerted a small influence on cementite precipitation, the cementite should also precipitate with the same Mn concentration as the bulk composition, while keeping the tie-line C₂D₂; in this case, the iso-C activity line at B₂; a_C,2^γ/α was obviously higher than that at C₂; a_C,2^γ/θ, which allowed the formation of the non-partitioning pearlite transformation. However, the C flux J_C,₂ was obviously lower than J_C,₁, because a_C,2^γ/α was lower than a_C,1^γ/α. Furthermore, if the bulk Mn composition is higher than the PLE/NPLE boundary for γ-α (c), ferrite with an optional composition A₃ may nucleate and grow under PLE mode. Because the proeutectoid ferrite does not form in this type of steel (Fig. 3(e)), cementite with a D₃ composition should precipitate, and the line A₃D₃ must pass through the bulk composition. At this time, the concentration in austenite at the ferrite/austenite and the cementite/austenite interfaces is denoted as B₃ and C₃, respectively. To promote the pearlite transformation, the iso-C activity line at B₃; a_C,3^γ/α must be higher than that at C₃; a_C,3^γ/α, while making the Mn partitioning between ferrite and cementite as small as possible. Under this condition, J_C,₃ must be very low. Thus, it can be concluded that the C flux controlling the pearlite transformation kinetics has the following order: (a) J_C,₁ > (b) J_C,₂ >> (c) J_C,₃. This corresponds to the experimental results in the pearlite transformation rate shown in Fig. 7. Conclusively, it should be noted here that the partitioning pearlite transformation is accompanied with Mn partitioning between the ferrite (A₃) and the cementite (D₃). Figure 9 shows the bright field STEM images exhibiting a growing pearlite interface, and the respective line profiles of Mn concentration across the ferrite/cementite and the ferrite/austenite interfaces in (a) 40C2Mn and (b) 40C3Mn. The measured line for Mn concentration profile corresponds to the white lines in each STEM image. Additionally, the bulk Mn composition and the position of the interfaces are indicated by the dotted and dashed lines in each line profile, respectively. In 40C2Mn, it was clearly observed (a) that Mn is uniformly distributed through the pearlitic ferrite and cementite (a-1), although the Mn-spike indicated by the black arrow exists at the ferrite/austenite interface (a-2). This certainly corresponds to the case shown in Fig. 8(b). Moreover, in 40C3Mn (b), Mn was clearly partitioned between the pearlitic ferrite and the cementite (b-1), and the Mn enriched region was confirmed in the austenite at the ferrite/austenite interface (b-2). This Mn partitioning behavior is in good agreement with the case shown in Fig. 8(c), which proves that an extremely slow transformation rate existed in 40C3Mn, and was controlled by Mn partitioning that occurred between the ferrite and the cementite. That is, LE theory can explain the transition between non-partitioning and partitioning pearlite transformations. In fact, the thermodynamic calculation using Thermo-Calc. (TCFE9 database) revealed that the C activities in the austenite at the ferrite/austenite and the cementite/austenite interfaces, which can be evaluated from the average Mn concentration in the ferrite (2.4%Mn) and cementite (7.0%Mn), respectively, were in close proximity to each other.

Fig. 6.

Schematic illustration of carbon flux controlling pearlite transformation.

Fig. 7.

Relationship between growth rate; v_P and interlamellar spacing; λ during pearlite transformation in Fe–C and Fe–C–Mn alloys.

Fig. 8.

Schematic isothermal phase diagram of Fe–C–Mn system explaining the variation of carbon activity gap controlling pearlite transformation kinetics depending on bulk Mn composition.

Fig. 9.

Bright field TEM images and Mn concentration profile across ferrite/cementite and ferrite/austenite interfaces in (a) 40C2Mn and (b) 40C3Mn.

It is also interesting that, in (a-1), Mn appeared to be locally depleted in the pearlitic ferrite at the ferrite/cementite interface in the 40C2Mn, as indicated by the arrow in (a-1). This Mn depletion zone was also observed in 50C2Mn and 60C2Mn. As shown in Fig. 8(b), the ferrite and cementite grew with the same Mn concentration as the bulk Mn composition when the non-partitioning pearlite transformation occurs. However, according to LE theory, thermal equilibrium should be realized between the ferrite and the cementite, even in the newly formed pearlite. Therefore, it is thought that Mn is redistributed between the ferrite and the cementite immediately after the pearlite transformation, which results in local partitioning between the ferrite and the cementite with the Mn depletion zone [4, 8]. Thereby, the thermal stability of the lamellar cementite in the pearlite may be markedly different between 40C0Mn and 40C2Mn (Figs. 3(b), 3(d)).

3.2. Effect of Mn Addition on Mechanical Property

3.2.1. Variation of Tensile Properties of Pearlite Steels by Mn Addition

Figure 10 shows the nominal stress-strain curves of all steels that were isothermally transformed at 873 K for 600 s, although the holding time for 40C3Mn was prolonged to 10.8 ks to complete partitioning pearlite transformation. The 40C0Mn exhibited moderate strain hardening accompanied with a clear yielding point, owing to the existence of a certain amount of proeutectoid ferrite. On the contrary, other steels had larger strain hardening, which led to a high tensile strength and sufficient total elongation, which is characteristics of pearlite steels. The tensile properties are listed in Table 2. Each value represents the mean evaluated by testing for three times. It was found that all steels had good tensile strength and a total elongation balance (σ_UTS × ε_total). However, there was a clear difference in the percent reduction of area at the fracture amongst them. Therefore, to discuss the effects of Mn on the strength and ductility in more detail, 0.2% proof stress (σ_0.2) and percent reduction of area (RA) balance are summarized in Fig. 11. By carrying out a comparison between the steels with and without Mn, it was found that σ_0.2 effectively improved by the Mn addition, regardless of the difference in the transformation mode between non-partitioning and partitioning pearlite. However, the effect of the Mn addition on RA can be classified into two types depending on the C content. When C content was lower than 0.6%, as indicated by the open circles, the Mn bearing steels maintained high RA with a higher σ_0.2 in comparison with the 40C0Mn. Moreover, the Mn addition caused a significant degradation of RA in the 100C steels, as indicated by the solid circles. Because the fracture surface was identified as an intergranular fracture in the 100C0Mn and 100C2Mn, the degradation of RA was caused by the thickening of the proeutectoid cementite that decorated the prior austenite grain boundaries. Furthermore, the retaining of RA in steels with a lower C content is attributed to the thermal stabilization of the lamellar cementite. Teshima et al.²⁴⁾ compared the local deformation and fracture behavior between two pearlite steels with lamellar and spheroidized cementite using a digital image correlation technique. Thus, they demonstrated that the fine lamellar structure of the cementite restricts the plastic deformation of the ferrite matrix, which leads to the enhancement of strain hardenability and the retardation of early fracturing. Therefore, it can be said that the effect of the Mn addition on the thermal stability of the lamellar cementite contributes to the development of local deformability in the pearlite.

Fig. 10.

Nominal stress-strain curves of steels isothermally transformed at 873 K.

Table 2. Tensile properties of steels isothermally transformed at 873 K.

	Isothermal holding	0.2% proof stress, σ0.2/MPa	Tensile strength, σUTS/MPa	Total elongation, εtotal/%	σUTS×εtotal/ MPa%	Percent reduction of area, RA/%
40C0Mn	873 K – 600 s	310	546	45.2	24700	62.9
40C2Mn	873 K – 600 s	370	727	37.3	27100	74.8
40C3Mn	873 K – 10.8 ks	393	749	34.2	25600	56.9
50C2Mn	873 K – 600 s	446	883	33.4	29500	60.8
60C2Mn	873 K – 600 s	436	943	32.3	30500	51.5
100C0Mn	873 K – 600 s	481	899	26.0	23400	20.3
100C2Mn	873 K – 600 s	515	1212	22.2	26900	9.0

Fig. 11.

0.2% proof stress and percentage reduction of area in steels isothermally transformed at 873 K.

3.2.2. Strength Mechanism of Pearlite by Mn Addition

There exist two possible mechanisms for the strengthening of pearlite steel. It is well known that the strength of the pearlite steel is determined by the interlamellar spacing of the lamellar structure. Marder and Bramfitt²⁵⁾ investigated which microstructural unit dominates the mechanical properties of pearlite steel, and demonstrated that its strength is solely dependent on the interlamellar spacing of cementite, and independent of the prior austenite grain size or block diameter. In particular, they emphasized that 0.2% proof stress and fracture strength are inversely proportional to the interlamellar spacing. Moreover, the authors recently reported that 0.2% proof stress linear increase with respect to the lattice strain (micro-strain) in the pearlitic ferrite, as evaluated by X-ray diffractometry, regardless of the interlamellar spacing and cementite morphology.²⁶⁾ Thus, authors proposed that the internal stress generated by the misfit between ferrite and cementite has large influence on the strength of pearlite steels.^{26,27,28,29,30,31)} Figure 12 shows X-ray (211) ferrite diffraction peak containing Kα1 and Kα2 of 40C0Mn and 100C2Mn with IF steel as a reference. As shown above in Table 2, 40C0Mn and 100C2Mn have the lowest and highest 0.2% proof stress, respectively. Indeed, comparing these diffraction peak, it is confirmed that pearlite steels have wider diffraction peak comparted with IF steel and furthermore the peak of 100C2Mn looks wider than that of 40C0Mn. Therefore, the 0.2% proof stress σ_0.2, which was evaluated by the tensile tensing, was plotted as a function of the inverse of the interlamellar spacing λ⁻¹ and lattice strain ε_L in the pearlitic ferrite, as shown in Figs. 13 and 14, respectively. Because the lamellar structure was broken out in 40C0Mn, its interlamellar spacing was temporally evaluated by the observation of the lamellar structure formed just after the pearlite transformation (Fig. 4). Additionally, the lattice strain in 40C0Mn reflects the internal stress not only in the pearlitic ferrite, but also in the proeutectoid ferrite, because this steel contains a certain amount of proeutectoid ferrite. Therefore, the 0.2% proof stress in the 40C0Mn may not depend on these two factors. Figure 13 shows that the interlamellar spacing in the Mn bearing steels was finer than that in the Mn free steels. Whether the Mn addition makes the lamellar structure fine or not has been investigated by previous studies.^5,7,9) However, Tashiro and Sato investigated the effect of many alloying elements on the pearlite structure comprehensively, and revealed that Mn slightly decreased the interlamellar spacing, possibly owing to the decrease in the interfacial energy between ferrite and cementite. Moreover, as can be seen in Fig. 14, the ferrite lattice strain increased by the Mn addition. Maejima et al.³²⁾ have investigated the effect of the microalloyed V on the microstructure and the mechanical property in eutectoid steel with 0.1% V using transmission electron microscopy and 3D atom probe microscopy. They concluded that the microalloyed V influences not the precipitation of VC in the lamellar ferrite, but the lattice strain increment in the pearlite lamellar through the partitioning of the solute V between ferrite and cementite, which led to an increase of 160–170 MPa in the 0.2% proof stress. According to their study, the Mn partitioning between ferrite and cementite is thought to exert the same effect as V. By comparing both figures, it can be observed that the liner relationship was well established both from the viewpoint of λ⁻¹ and also from the viewpoint of ε_L. Therefore, it is difficult to determine the dominant strength mechanism for the pearlite steel. Although further investigation is required on the strengthening mechanism of pearlite, it can be concluded that the Mn addition improved the strength and local deformability balance in the pearlite steel, unless the proeutectoid cementite precipitated at the previous austenite grain boundaries.

Fig. 12.

X-ray (211) ferrite diffraction peak of IF steel, 40C0Mn, and 100C2Mn.

Fig. 13.

Relationship between 0.2% proof stress and interlamellar spacing in steels isothermally transformed at 873 K.

Fig. 14.

Relationship between 0.2% proof stress and lattice strain of pearlitic ferrite in steels isothermally transformed at 873 K.

4. Conclusions

The effects of the Mn addition on the microstructure formed via isothermal transformation at 873 K and its tensile properties were investigated in medium- and high-carbon steels with 0.4–1.0 mass% C. The obtained results are as follows:

(1) Mn inhibits the formation of proeutectoid ferrite. This is attributed to the retardation of the ferrite transformation rate. Additionally, the pearlite transformation rate was retarded by the Mn addition, which led to the relatively slow non-partitioning transformation and an extremely slow partitioning pearlite transformation. Particularly, the transformation rate of the partitioning pearlite transformation associated with the Mn partitioning between ferrite and cementite at a transformation interface was more than three orders of magnitude lower than the pearlite transformation without Mn at a given interlamellar spacing.

(2) The retardation of pearlite transformation by the Mn addition is caused by the decrease in carbon flux at the transformation interface controlling the pearlite transformation, which can be explained on the basis of local equilibrium theory at austenite/ferrite and austenite/cementite interphases. Assuming that local equilibrium theory is satisfied even at ferrite/cementite interface in the pearlite, Mn partitioning between ferrite and cementite should occur immediately after the pearlite transformation. This causes the thermal stabilization of lamellar cementite.

(3) The Mn addition caused the refinement of interlamellar spacing and the increase of lattice strain in pearlitic ferrite. One or both of them contributed to the strengthening of pearlite steel. Consequently, tensile properties, particularly the strength and local deformability balance, were developed by the Mn addition in pearlite steel, regardless of the difference in the transformation mode between the non-partitioning and partitioning pearlite, unless the proeutectoid cementite precipitated at the previous austenite grain boundaries.

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